Heat treatable beta titanium-base alloys and processing thereof



Sept. 1, 1 M. B. VORDAHL 3,147,115

HEAT TREATABLE BETA TITANIUM-BASE ALLOYS AND PROCESSING THEREOF Filed Sept. 9, 1958 3 Sheets-Sheet 2 Fl /cm. QZ OFFSET YIELD 577?ENGTH- wE/aHr COMPA/P/SON F AGED HIGH .STEENGrH SHEET ,47' 5(ELONG'AT/0N LEVEL 5/20 VCA 7743.51 l/C'r' -3,4Z

SUPEEALLO) j TEMPEEWTUIP- F 0 INVENTOR. ML TONE V PQAHL NG THEREOF Sept. 1, 1964 M. B. VORDAHL HEAT TREATABLE BETA TITANIUM-BASE ALLOYS AND PROCESSI 3 Sheets-Sheet I5 Filed Sept. 9, 1958 I'VE/G711 COMPARISON E C/PEEP STRENGTH- OF #650 SHEEr MAT R/ALS BASED ON STRESS 7'0 PRODUCE QZZPmSnc C/PEEP //v .700 Hou/PS CbMPos/r/o/v M5? VG DO C /20,4 V 77 5-7 /5 /7-7 I7 O 7 EMPElP/7 rue: F

INVENTOR. ML TON 5 VOEDAHL. BY

United States Patent 3,147,115 HEAT TREATABLE BETA TITANIUM-BASE ALLOYS AND PROCESSING THEREOF Milton B. Vordahl, Beaver, Pa., assignor to Crucible Steel Company of America, Pittsburgh, Pa., a corporation of New Jersey Filed Sept. 9, 1958, Ser. No. 759,973 8 Claims. (Cl. 75-175.5)

This invention pertains to heat treatable titanium-base alloys and to methods of heat treating the same.

The invention pertains more particularly to titaniumbase alloys consisting essentially of about 10 to 16% vanadium, 9 to 13% chromium, 2 to 5% aluminum, which are heat treatable by solution annealing in the temperature range extending from slightly below to above the beta transus, preferably at about 12001600 F., and air cooling to impart a substantially all-beta structure and a relatively low yield strength for ease of forming, and which may thereafter be age-hardened on prolonged heating at about 800-ll00 F. to extremely high strength levels with retention of adequate ductility for service.

The alloys of the invention thus contain, as essential alloying ingredients, the isomorphous beta promoter vanadium and the sluggishly eutectoid beta promoter chromium, together with the alpha promoter aluminum. Part of the chromium may be substituted by manganese up to about 8%, or by iron up to about 4% on an equal weight percent basis, these elements like chromium being beta promoters of the sluggishly eutectoid type. Also part of the alumium may be substituted by up to about 4% tin, on a three-to-one weight percent basis, since three percent tin is equivalent to about 1% aluminum as regards its strengthening action and effect on ductility. The interstitial content of these alloys, i.e., carbon, oxygen and nitrogen, should preferably not exceed about 0.02% in total amount.

Preferred alloys of the invention are those containing about 12-14% vanadium, -12% chromium and 34% aluminum, excellent compositions being the following, Ti-l3V-1 1Cr4Al and Ti-13.5Vl 1Cr-3Al.

The alloys of the nivention possess a truly remarkable combination of properties and are the most versatile titanium-base alloys thus far developed. For example, they have the highest strength-weight ratio of any structural material developed to date, and the highest strength and best formability of any titanium-base material heretofore known. In addition they are ductile-weldable, cold-headable and have very great and deep hardenability, which fits them ideally for such applications as high strength-weight honeycomb foil at one extreme and for very thick heat treat sections at the other.

The high beta content of these alloys stabilizes titaniums high temperature beta structure down to room temperature in order to impart the advantages of the excellent formability of the beta structure, and also to render the alloys ageable to high strength levels at conveniently low temperatures. Addition of the alpha promoter, aluminum, or aluminum and tin, imparts added strength and provides the proper balance of the alloying elements so that the structure stays all-beta during slow cooling from the solution annealing temperature and dur- 3,147,115 Patented Sept. 1, 1964 ing forming, and permits of aging to quite high strength levels in conveniently short times at low temperatures, but without over-aging to a brittle or weak condition in service.

Although so-called formageable Ti-base alloys of the mixed alpha-beta type have previously been known which may be solution treated and quenched to a low yield strength for ease of forming and, in this condition, are susceptible to substantial age hardening, such alloys must be Water quenched from the solution treating temperature and are usually only marginally ductile in welded sections. In marked contrast, the formageable beta alloys of the instant invention need be only air cooled from the solution treating temperature and are ductile-weldable.

The beta alloys of the present invention, like their austenitic steel counterpart (188 stainless steel), have about one-quarter of their content as alloying elements, and this have a strong solid solution base and great strengthening potential, while being quite formable in the solution treated and air cooled condition. In this formable condition, the alloy has, for example, a minimum yield strength of about 120,000 p.s.i. Aging the formable alloy for about 24-96 hours in the 800-850 F. range strengthens it to about 190,000-240,000 p.s.i. in ultimate strength and to aboutl70,000220,000 p.s.i. in 0.2% olfset yield strength with good resulting ductilities. Considerably higher strength values of up to 280,000 p.s.i. in ultimate strength and 270,000 p.s.i. in yield strength, are obtainable with some reduction in ductility.

A typical alloy of the invention when aged to a tensile elongation level of 5% has a 0.2% oiiset yield strength of 190,000 p.s.i., which compares with a corresponding value of about 160,000 p.s.i. for the formageable alphabeta alloys above referred to and a value of about 220,000 p.s.i. for the PH steels, which values when converted to the strength-Weight ratio basis give a ratio of about 1,300,000 for the alloy of the instant invention as compared to about 975,000 for the formageable alpha-beta alloys and about 760,000 for the PH steels. I am aware of no currently available structural material having as high strength-Weight values as the alloy of this invention.

In the annexed drawings:

FIGS. 1 and 2 are constitution or phase diagrams of typical alloys according to the invention, in which, in FIG. 1, aluminum is a variable element, and, in FIG. 2, chromium is a variable element.

FIG. 3 is a graphical comparison on a strength-weight ratio basis, at temperatures extending up to 1000 F., of the 0.2% offset yield strength of aged alloys according to the present invention versus previously known competitive materials in the aged condition, including other age hardening titanium-base alloys, steels and superalloys.

FIG. 4 is a comparision, analogous to that of FIG. 3, except that in FIG. 4 the comparison is on a creep strength-weight ratio basis of the same materials as in FIG. 3. t

The following Table I gives tensile tests on various alloys according to the invention as solution treated and aged for the times and at the temperatures indicated, the

o the data being typical of the properties obtainable by the simple aging heat treatments shown.

TABLE I The hardening effect of low temperature aging tends to disappear at higher aging temperatures. The fine Longitudinal Tensile Properties of Simple Aged Sheet, Plate, Bar and Billets Sol. treat. Aging-treat. Tensile properties, p.s.i. 1,000 Composition, percent, Product bal. Ti

Hrs. F. Time, Temp, UTS 2% Percent Percent hrs. F. YS E1 R 13V-11Cr-3A1 1, 400 72 850 218 201 5.0 13. 13V-11Cr-3 kl 1, 350 64 900 202 184 5.0 3. 8 13V-11Cr-5Al 1, 400 64 800 206 100 3. 3 3. 9 13V-11Cr-4A1 1 1 1, 400 32 800 162 154 3.0 7. 13V-11Cr-3A1 1 1, 400 16 900 192 178 4. 0 6. 4 13V-11Cr-2%Al 1 1, 400 32 850 187 176 1. 5 3. 9 13V11Cr-2 ,Al 1 1, 450 64 850 205 204 1. 0 2.1 13V11Cr3A1. 1% 1,400 64 900 205 193 1.0 1. 0 13V-110r-4Al 2 1, 500 32 900 203 198 1. 0 1. 6 13V-11Cr-2%r 2 1, 450 G4 850 213 210 1. 0 2. 0 13V-11Cr2%Al 2 1, 400 850 181 169 3. 5 5. 4 13V-11Cr-5A1 2 1, 400 48 800 203 192 2. 0 2. 5

All specimens air cooled from the solution treating temperature.

Referring to the constitution diagrams for these alloys shown in the annexed FIGS. 1 and 2, it will be seen that the solution treating temperatures employed in the Table I tests were somewhat above the equilibrium beta transformation, although temperatures slightly above or even somewhat below the equilibrium temperature may be used. Examination of the constitution diagrams of these figures reveals that these are unstable beta alloys which are strengthened by transformation of some beta to alpha and by precipitation of T iCr compound after long time aging at about 800950 F. This precipitation phenomenon involving both phases is quite complex in nature but very advantageous in producing highest strength and ductility combinations by various aging treatments as described below. For the alloys containing manganese or iron in substitution for part of the chromium as aforesaid, titanium-manganese or titanium-iron compounds will of course be precipitated along the TiCr to impart strengthening action during aging.

The Table I tests establish that relatively slow cooling, such as air cooling, of these beta alloys is adequate for retention of ageability. This is a tremendous practical advantage over the formageable alpha-beta alloys both in mill productivity (solution quenching) and in shop fabricability, e.g., brazing and in quenching shapes or thin sections.

My tests have shown that cold or warm working the beta, i.e., working below the recrystallization temperature, produces a uniform distribution of the subsequently transformed alpha and TiCr or equivalent Ti-Mn or Ti-Fe compounds, and also accelerates the rate of hardening on aging.

The strengthening due to age hardening probably involves omega formation in addition to compound formation and coherent alpha separation. The hardening due to alpha separation is good as when nucleated or distributed by warm or cold working, supra, but may become harmful to ductility if preferentially present as networks or on preferred planes. Compound formation is likewise good if similarly well distributed by cold or warm working. Omega formation, however, appears undesirable. It decomposes rapidly at about 1000 F. and above, but slowly below this temperature range. Thus low temperature aging must be prolonged to eliminate it.

alpha and compound precipitated at about 800 F. redissolves at about 1000-1100 F. and a new hardening cycle is then initiated. Omega disappears very rapidly at 1000ll00 F. Alpha separated at high temperatures of about 1200 F. up to the beta transus, is apparently massive and discontinuous enough, that its effect on properties is slight. The remaining beta is enriched, and omega formation in subsequent aging inhibited. Aging in the range of 1000-1100 P. if prolonged tends to impart lowered ductility, while short time aging produces less continuous alpha with some benefits as above noted.

Based on the foregoing considerations, I find that aging temperatures between 800 and 950 F. appears to be optimum with aging times varying from about 24 to 96 hours. In general, long times at low temperatures produce the best combinations of very high strength and ductility. Slow decomposition of omega is believed to be responsible for this behavior. However, the necessity of longer times of to 500 hours at 650750 F. aging temperatures make this heat treating process commercially impractical.

Accordingly, the best results as regards high strength combined with good ductility based on commercially practical aging time, are not obtainable with the simple aging treatment of Table I, and my investigations have shown that improvements may be effected by using double aging heat treatments comprising either a low temperature aging preceding a high temperature treatment or vice-versa. As to the former, I have found that a retrogression or reversion occurs on heating low temperature, aged alloys according to the invention for short times at higher temperatures. A new cycle roughly characteristic of the higher temperature then starts. This may be due to compound resolution, but the exact nature of change is unknown. Employment of this method of aging, however, has been found to improve the percent elongation and area reduction of these alloys with slight sacrifice in strength. Experiments covering this procedure establish that aging at about 48 to 96 hours at temperatures of about 800 to 900 F., followed by further aging for about 10 minutes at about 1000-1050 F. produces optimum results. Typical examples of this are shown in the following Table II.

TABLE 11 Longitudinal Tensile Properties of Double Aged (Low Temperature Preceding High) Sheet, Plate, Bar and Billet I Sol. treat. Tensile properties, p.s.i. X 1000 Oompositlon, Product percent, bal. Ti Aging treatment Hrs. F.* UTS 2% Percent Percent YS E1 RA 13V-11Cr-3Al.. 1,400 48 hr. at 900 F.+2 185 169 10. 15. 9

min. at 1100 F. 13V--11Or--3Al 1,400 48 hr. at 900 F.+1 184 106 8. 3 13.0

min. at 1100 F. 13V-11Cr-5A1 1 1, 400 64 hr. at 900 F.+30 228 213 3.3 3. 5 min. at 100 13V11Cr-4A1 1 1, 400 64 hr at 900 F +15 228 220 2.0 1.6

mm at 1050 F 13V-11Cr3Al 1 1, 400 64 hr at 900 F +1 187 176 5. 0 8. 6

hr at 1050 F 13V110r2%Al 1 1, 450 64 hr at 850 F 195 186 2. 0 3. 0 min. at 1000 F. 13V11Cr3Al 1% 1, 400 64 hr. at 900 F.+15 202 192 2. 5 2. 1 min. at 1050 F. 13V110r2%A1- 2 1,400 64 hr. at 850 F.+1 185 177 3.0 7.2

hr. at 1000 F. 13V11Cr5A1 2 1, 400 200 hr. at 900 F.+1 228 226 1.0 0.8

hr. at 1050 F.

* All specimens air cooled.

Comparing the test results of the double aging treatment of Table II with the simple aging treatment of Table I, it will be seen that there is a noticeable improvement in tensile elongartion and area reduction at very little sacrifice in strength.

A double aging treatment employing a high temperature aging preceding a low temperature aging appears to involve separation and agglomeration of alpha and enrichment of beta above the compound forming temperature, with subsequent strengthening by both further alpha separation and compound formation at a lower temperature corresponding to the three-phase field in the FIG. 2 phase diagram. The advantage of using short time, high temperature aging before a long time, low temperature treatment is to avoid unfavorable alpha distribution and to minimize initial omega formation. Some typical results obtained by using this aging procedure are given in the following Table III, which show fairly good strengthductility combinations.

The test results of Table III illustrate what I have found to be generally true of these alloys, namely, that in general the ductility remaining after aging is less with heavy sections than with light or thin gauge stock. That is to say, the properties of aged sheet material are usually better than with bar and billet stock. However, since the latter are more difficult to cold or warm work than is sheet material, the advantages of the double aging heat treatment are of greater importance and needed more with heavier sections, where applicable, than with sheet.

TABLE III The beneficial eifects on the aged properties of warm rolling in the alpha-beta temperature field prior to aging, i.e., rolling below the recrystallization temperature, of about 1400 F., are shown by the test results in the following Tables IV, V and VI.

TABLE IV Efiect of 10, 30, Warm Rolling at 1200 F. on Heat Treated Properties of /8 Sheet Bar Heat number: R4942 (13V-11Gr-4Al) Specimen type: .150 diameter, standard tensile Processing sequence: Solution treatment+warrn roll+age Percent 2% Per- Perreduetio Aging treatment UTS YS cent cent at F. El BA 10 at 1200.. V 147 141 18.3 36.2 30 at 1200.. 4 hrs. at 800 F.AC 150 143 16. 7 36. 2 60 at 1200.. 161 152 16. 7 33. 9

10 at 1200 168 159 13. 3 29. 7 30 at 1200.. 16 hrs. at 800 F.-AC 181 169 13. 3 26. 6 60 at 1200 195 183 10.0 23.5

10 at 1200.. 234 220 1. 7 4.0 30 at 1200.. 64 hrs. at 800 F.AO 235 225 3. 3 7. 6 60 at 1200.. 247 236 3.3 6.4

10 at 1200-. 64 hrS. at 900 F.+10 min. 228 209 3. 3 4. 7 20 at 1200.- at 1000 F. 232 216 3.3 6. 3 60 at 1200.. 247 230 3. 3 5. 2

Longitudinal Tensile Properties of Double Aged (High Temperature Preceding Low) Sheet, Plate, Bar and Billet S01. treat. Product Composition, Aging treatment UTS 2% Percent Percent percent, bal. Ti YS E1 RA Hrs. F.*

Sheet 13V11Cr-3Al. 1 1, 400 2 hr. at 1200 F.+ 185 167 5.0 7. 6

48 hr. at 900 F.+ 2 min. at 1100 F. Do 13V11Cr3A1 1,400 4 hr. at 1200 F.+ 191 6. 7 8.0

72 hr. at 900 F. Plate 13V11Cr-5Al 1 1, 400 16 hr. at 1400 F.+ 194 177 3. 3 2. 8

' 48 hr. at 900 F. Bar 13V-11Cr-5A1 1 1,400 %hr. at 1200 F.+ 194 172 2.0 3. 3

64 hr. at 900 F. Do 13V-110r-4A1 2 1, 500 4 hr. at 1200 F.+ 176 1.0 1.0

64 hr. at 900 F. Billet 13V-11Cr-5A1 2 1, 400 4 hr. at 1200 F.+ .183 171 2.0 2. 4

64 hr. at 900 F.

"All specimens air cooled.

TABLE VII TABLE VIII 30 and 40% Cold Rolling on Heat .100 thick microteusiles Heat number: P2-13060l (13V-11Gr-3Al) Specimens: Standard flats Heat Treat Data on .010" Face Sheet Cold Rolled About 70% After Annealing at 1400 F.

Heat number: R4042 (13V-11Cr-4Al) Specimen type: Processing sequence: 1400 F. ST+cld roll-l-age 0018 7978 2 @MM wwmm EH M P -t 7730 7037 3 u m n ns? 01 8 PC www Mu m W 1122 22%% 2 2155 2 6 0905 6 S T mama amam aaaa a U 0 n u u n H m o o v o e .A A A A A m .R m R F F F F0 8 0 0 0 0 m m m a my M a a M Ma s s s sn S t r l A r h h H .mm h 6 4 4. 4 4 1 no 6 6 i1 1 i i Treated Properties of .1l0.145T hick Sheet prior to aging is shown by the test results in the following Tables VII and VIII.

standard tensile treatment-t-rvarm rolH-age 7 TABLE V Treated Properties of Sheet Bar diameter,

Heat number: R4942 (13V-11Gr-4A1) gpecimen type: .150

rocessing sequence: Solution Eflect of 10, 30, 60% Warm Rolling at 1300 F. on Heat on ,m mmwa i 11 1. 7 9 mmflfl 2H2 p t na fl fi 210 21 0 PB m t W GMM 4.01 8553 B 567 .7759 Oma -W W 222 2222 t S em .lmng 160 I068 Hum 667 7. 9 222 0 -2 2 t S t. 1 H 0 e r 200 2000 m mh 7240; t T 12n t r v. t e g m n .m tr w m mmm mama mm w 6 6 777 7777 e A p 0 0 5 0 1 2 2 3 167 706 692 A r w Q O 877 579 e AWE Pc h c t R 1 303 333 307 M W HmE 33 3 35 6 a w PC 2 0 0 r. 7s 1 I t m BY 026 455 569 4 a 0 0 444 2 (1 W 111 mzfi W 00 s rm 0 T al DA U s when mow u e mam B 0 T 111 111 2%2 222 I t R F r U V w y 0% i t t b m u u u E S n t u n L 0 V m l n u n 1 B 1 3 n t C C m 0 C m 0 m C A A m M a u 21 C m A I' W: t F F a 8 a F F R n m 0 m m 0 D Re t. 0 8 8 %F 1% ms 8 t t to n M .1 M a a a0 8 4 bm t E s s w m 8 t n A s m m r 1. :1 n m m 6 4 we a 0 mm t 4 l 6 0 N t M 0 l ll ill it, c 1 m r T e 00 t. n a HS t Ht rR 00m mmo 00m 0mm m E but 333 333 333 333 mw v 111 111 111 111 e H i new a e mam Mam H M 000 00 000 000 36 36 136 136 *Broken into several pieces, it was impossible to measure the elongation. 35

Efiect of 10, 20

Percent cold reduction at 800 F.+ min at 1050 F. 64 hrs m; 800 F.+

min at 10501.

These tests clearly establish that small reductions by warm rolling at about 1200-1400 F. prior to aging, is

8 hrs. at 800 F.AC. As received.-.

32 hrs 64 hrs 64 hrs very beneficial in improving the ductility without appre- 10- ciable sacrifice in strength. The warm rolling appears 5O to cause compound or alpha nuclei to form at lower tem- :11-"--. peraturc during aging, and also to be more uniformly distributed than without such working. As shown by the 10mm test results, a single aging under these circumstances pro- 40 duces sufiiciently good results that it is not essential to ernploy double aging after warm rolling. From Tables 1V and V, inc., it will be seen that tensile strengths up to approximately 200,000 p.s.i. are obtained in the aged Referring to Table VIII, it will be seen Specimen broke in grips. strengths are obtainable with the previously warm rolled 35 It will be seen from these tables that by cold rolling prior to aging, it is possible to get ultimate strengths of 250, GOO-300,000 p.s.i. with accompanying tensile elongations of about 0.5 to 2.5%, which on a strength-weight ratio basis would correspond to about 500,000 p.s.i. in a sheet to sheets and relatively light gauge plates made of alloys according to the invention, it thus shows great promise in producing very high strengths with desirable ductilities due to favorable distribution and orientation of the hardspecimens along with tensile elongations of about 10% and 10--- area reduction of about 24% as compared to much lower 20 such values obtained with the aging procedures of Tables 30: I to III for the same tensile strength. It will further be noted from a comparison of Tables I-III, inc., with those of Tables IV-VI, inc., that in general much higher aged alloys than where this step is omitted, the Warm rolled, aged strengths ranging up to about 250,000 p.s.i. with corresponding tensile elongations of about 3% and area reductions of about 6%, as compared to tensile strengths ranging up to only about 230,000 p.s.i. with accompanymade of steel. Although cold rolling can be applied only ing elongation and area reduction values of only about 1% when the step of warm rolling is omitted. Thus, the warm rolling introduces a vast improvement in the properties of the subsequently aged alloys.

The efiect of cold work on the alloys of the invention ening phases.

that the rate of aging increases with the amount of cold work, so that for a given strength level the use of shorter aging times on cold rolled material may be employed than where this step is omitted, thus rendering the aging operation more efllcient and economical. Thus, reverting to Table I, it will be seen that a 64 hour aging cycle was required to obtain an ultimate strength level of about 200,000 p.s.i., whereas as shown in Table VIII, this strength level can be obtained in as little as 4 hours where the material has been subjected to a 40% cold reduction prior to aging. Another vitally important improvement to be noted with reference to the aged properties of the previously cold rolled material, is the excellent ductility after aging as shown both by percent tensile elongation and percent area of reduction. For example, in Table VIII, at an ultimate strength level of 200,000 p.s.i., these previously cold rolled alloys showed elongations of 10% and area reductions of about 30%, as compared, reverting to Table I, to percent elongation and area reduction values of about 3-4% for an alloy aged to the same strength level but without prior cold rolling. Also from a comparison of the warm rolling test results of Tables IV-VI with the cold rolling test results of Tables VII and VIII, it will be seen that the cold rolling prior to aging results in a much higher aged ultimate strength at a given ductility level than is obtained with the previously warm rolled materials, the aged strength in both instances, however, being higher than those obtainable where the warm or cold rolling is omitted, again at a given ductility level, as shown by comparison with the test results of Tables IIII, inc.

However, the discovery that warm work, as in Tables IV-VL inc., accomplishes much of what cold work does, as in Tables VII and VIII, is of particular importance with sizes and shapes which would be difficult to cold work, yet relatively easy to warm work. For example, a heavy dished or flanged end of a tank could be warm spun with relative case, but not cold drawn or spun, etc.

I have found that in welded articles made of the alloys of the invention, the weld metal and the heat affected weld zone age more rapidly than the base material at any given aged strength, and that at any given aged strength, the weld is less ductile than the base material, unless precautionary measures are taken to overcome these objectionable features. One way to minimize these effects is to partially age and then welded and then age a little more. In this way the base gets a lot of aging and the weld gets a little aging and the strength throughout is reasonably close. Another procedure is to cold or warm work the material prior to fabrication and Welding. As a result, since the weld has no cold work, of course, the cold or warm worked base material and the Weld, age at more or less the same rate to similar strengths and ductilities. Test results on welded structures employing these procedures are shown in the following Table IX.

TABLE IX Various Heat Treatment Processes to Obtain High Strength-Ductility Combinations n Welded Sheet and Thin Plate I have found that the hardening mechanism of alloys according to the invention, is quite suitable for any commercial brazing application such as honeycomb construction. For high strength, brazed structures, it is necessary to heat treat after brazing. Furnace cooling after brazing very conveniently produces a solution treated structure which is readily ageable to a higher strength. I am aware of no other titanium-base alloy which is sluggish enough to retain age hardenability after slow cooling from the brazing cycle. As a result of this metallurgical characteristic of the alloys of this invention, all the difficulties arising due to faster cooling rates (air or water) from solution temperature, i.e., distortion, warpage, handling problems, costly jigs, etc., are eliminated in complex brazing structures such as honeycomb constructions by employment of the alloys of this invention. There fore, from a practical and economical standpoint, these alloys are tailormade for brazing applications. The experimental results of a typical brazing and heat treating operation are given below, for an alloy of analysis of composition Ti-13V-1 lCr-4Al.

As compared to the above, in practice usually all other age hardening alloys have to be water quenched from a single phase field prior to aging, such as in the formageable titanium-base, alpha-beta alloys as Well as alloys of aluminum or copper base and the like.

Since yield strength is an important criterion as regards service performance of structual materials, FIG. 3 compares on a strength-weight basis at temperatures up to 1000 F., the 0.2% offset yield strengths of aged sheet materials made of the formageable beta alloys of this invention with the formageable alpha-beta alloys, the PH or precipitation hardening steels and with the superalloys. All specimens were aged to strength levels imparting a tensile elongation of 5%, at which the corresponding room temperature yield strengths obtained were as follows: PH steel, 220,000 p.s.i.; BIZOVCA, 190,000

p.s.i.; C120AV, 160,000 p.s.i.; and the superalloy A286,

70,000 p.s.i. Other formageable alpha-beta alloys tested gave the same performance graph as C120AV, namely C115AMoV (Ti4Al3Mn-IV) and CVA (Ti-16V-2.5Al).

It will be seen from FIG. 3, that compared on a strength-weight basis, the formageable Ti-beta alloys according to this invention are markedly superior in aged strength at all temperatures up to 1000 F., to the formageable Ti-alpha-beta alloys, and even more so as regards the PH steels and the superalloys. If the aged strength comparison is made at a tensile elongation level of 2%, the yield strength of the PH steels is increased to 260,000 p.s.i. and for the formageable beta alloys of the invention to upwards of 220,000 p.s.i. The corresponding strength-weight comparison thus becomes about 1,280,000 for the beta alloys versus about 925,000 for the PH steels. Accordingly, whether compared at the aged level of 5% or 2% tensile elongation, the strength of the aged beta alloys of the invention is appreciably better than that of the PH steels. Similar comparisons on an ultimate strength basis, show the same outstanding performance of the aged beta alloys of this invention as compared to the other structural materials above discussed. The above and similar tests, therefore, establish that the aged alloys of the present invention have the highest strength-weight value of the known structural sheet materials that are currently available.

Creep strength is another important criterion as regards service performance, with respect to which FIG. 4 gives a comparison analogous to that discussed with reference to FIG. 3, the graph designations being the same in each. In FIG. 4, as noted in the heading, the comparison is on the basis of the stress to produce 0.2% plastic strain in 500 hours, the sheet specimens employed again having been aged to strength levels corresponding to a tensile elongation level of 5%. The FIG. 4 comparison like that of FIG. 3 is on a strength-weight basis at temperatures up to 1000 F.

As thus compared the superiority of the Ti-beta alloys of this invention as regards creep strength is shown in FIG. 4 to be achieved at all temperatures up to at least 600 F. as compared to the formageable Ti-alpha-beta alloys, the PH steels and the superalloy, and the indications are, as shown by extrapolation of the Ti-beta curve, that this superiority exists up to 800 or 900 F., above which the PH steels and superalloy show superiority.

For the alloys of this invention in the form of solution treated, or annealed sheet, typical properties at various temperatures are given in Table X below.

Corresponding notched impact values are given in Table XI below.

TABLE XI Typical Notched Tensile Properties (Annealed Condition) Temp, Notched Unnotched Unnotehed Unnotched Unnotehed F. ultimate ultimate yield elong., RA

percent percent .005 R60 F. notch.

Creep and creep stability test results for the annealed alloys are given in Table XII below.

TABLE XII Creep and Creep Stability (Annealed Condition) Initial Final Total Stress, Test deiormadeformaplastic Temp, p.s.i. duration, tion, tion, deforma- F. hours percent percent tion,

percent 89 110, 000 1, 007 806 794 Nil 80 120, 000 959 S40 925 085 Typical mechanical properties of aged sheet material of these alloys is given in Table XIII below.

12 TABLE XIII Room Temperature Tensile Properties (Aged Condition) 0.2% offset Ultimate Percent yield, p.s.i. strength, p.s.i. elongation Elevated temperature properties of the aged sheet material are given in the following Table XIV.

TABLE XIV Elevated Temperature Tensile Properties (Aged Condition) Ultimate Temperature, F. strength, 0.2% offset Percent p.s.i. yield, p.s.i. elongation Notched impact values of the aged material are given in Table XV below.

*Determined on 5 round bar.

Creep strength and thermal stability of the aged material are shown in Table XVI below.

TABLE XVI Creep Strength and Thermal Stability (Aged Condition) I Room temperature Creep condition properties after creep exposure Stress, Temp., Time, Per- UTS, 0.2% Elong k.s.i. hrs cent k.s.i. yld., percreep l:.s.i. cent Heat1. 575 1, 435 10 211 196 100 575 1, 507 10 210 188 5 100 600 712 16 215 192 6% Heat 2 95 600 1, 502 11 200 184 1 Gage mark break. 2 Gage mark break with 2% elong; 15% RA The creep stresses required to produce the 0.2% plastic deformation in 500 hours shown for Ti-beta graph of HG. 4 are:

Temperature, F. 0.2% plastic deformation, p.s.i.

Charpy V-notch impact properties of the aged material are given in Table XVII below.

TABLE XVII Chai'py V-Notch Impact Properties (Aged Condition) Charpy impact value,

Temperature: foot-pounds This application is a continuation-in-part of my copending applications Serial No. 671,316, filed July 11, 1957 (now Patent 2,857,267); Serial No. 729,830, filed April 21, 1958 (now abandoned); Serial No. 549,164, filed November 25, 1955 (now abandoned); and Serial No. 435,754, filed June 10, 1954 (now Patent 2,974,076); which are in turn continuations-in-part of my applications Serial No. 229,143, filed May 31, 1951 (now Patent 2,950,191), Serial No. 132,327, filed December 10, 1949 (now abandoned), and Serial No. 356,877, filed May 22, 1953 (now Patent 2,754,203).

What is claimed is:

1. A heat treated and age hardened titanium-base alloy consisting essentially of about: to 16% vanadium, .9 to 13% chromium, 2 to 5% aluminum, up to 8% manganese, up to 4% each of iron and tin, up to 0.2% in total amount of carbon, oxygen and nitrogen, balance substantially titanium, characterized by an ultimate strength on the order of 190,000240,000 p.s.i., an 0.2% oifset yield strength on the order of 170,000-220,000 p.s.i. and a tensile elongation of at least 1%.

2. A heat treated, age hardened titanium-base alloy consisting essentially of about: 10 to 16% vanadium, 9 to 13% chromium, 2 to 5% aluminum, up to 0.2% in total amount of carbon, oxygen and nitrogen, and the balance substantially titanium, characterized by an ultimate strength on the order of 190,000240,000' p.s.i., an 0.2% ofiset yield strength on the order of 170,000- 220,000 p.s.i., and a tensile elongation of at least 1%.

3. The method of heat treating and fabricating a titanium-base alloy consisting essentially of about: 10 to 16% vanadium, 9 to 13% chromium, 2 to 5% aluminum, up to 8% manganese, up to 4% each of iron and tin, up to 0.2% in total amount of carbon, oxygen and nitrogen, balance substantially titanium, which comprises: solution annealing said alloy at a temperature extending from slightly below to above the beta transus, air cooling substantially to room temperature, forming said alloy to shape, and thereupon strengthening said alloy by aging at about 700 to 1100 F.

4. The method of processing a titanium-base alloy consisting essentially of about: 10 to 16% vanadium, 9 to 13% chromium, 2 to 5% aluminum, up to 8% manganese, up to 4% each of iron and tin, up to 0.2% in total amount of carbon, oxygen and nitrogen, balance substantially titanium, which comprises: solution annealing said alloy at a temperature extending from slightly below to above the beta transus and cooling to room temperature, thereupon plastically deforming the alloy at a temperature below the recrystallization temperature, and thereupon strengthening the alloy by aging at a temperature of about 700 to 1100 F.

5. The method of age hardening a titanium-base alloy consisting essentially of about: 10 to 16% vanadium, 9 to 13% chromium, 2 to 5% aluminum, up to 8% manganese, up to 4% each of iron and tin, up to 0.2% in total amount of carbon, oxygen, and nitrogen, balance substantially titanium, which consists in solution annealing said alloy at a temperature extending from slightly below to above the beta transus and cooling substantially to room temperature, thereupon aging for about 24 to 96 hours at a temperature of about 800 to 900 F., and thereupon further aging for about 1 to 30 minutes at a temperature of about 1000 to 1100 F.

6. The method of age hardening a titanium-base alloy consisting essentially of about: 10 to 16% vanadium, 9 to 13% chromium, 2 to 5% aluminum, up to 8% manganese, up to 4% each of iron and tin, up to 0.2% in total amount of carbon, oxygen and nitrogen, balance substantially titanium, which comprises solution annealing said alloy at a temperature extending from slightly below to above the beta transus and thereupon cooling substantially to room temperature, thereupon aging at a lower temperature in the range of 800 to 1100= F. for a relatively long time, and thereupon aging for a relatively short time at a higher temperature within said range.

7. The method of age hardening a titanium-base alloy consisting essentially of about: 10 to 16% vanadium, 9 to 13% chromium, 2 to 5% aluminum, up to 8% manganese, up to 4% each of iron and tin, up to 0.2% in total amount of carbon, oxygen and nitrogen, balance substantially titanium, which comprises solution annealing said alloy at a temperature in the range extending from slightly below to above the beta transus and cooling substantially to room temperature, thereupon aging at a higher temperature in the range of about 800 to 1200 F., and thereupon further aging at a lower temperature within said range.

8. The method of processing a titanium-base alloy consisting essentially of about: 10 to 16% vanadium, 9 to 13% chromium, 2 to 5% aluminum, up to 8% manganese, up to 4% each of iron and tin, up to 0.2% in total amount of carbon, oxygen and nitrogen, balance substantially titanium, which comprises solution annealing said alloy at a temperature extending from slightly below to above the beta transus and cooling to room temperature, thereupon cold Working said alloy, and thereupon strengthening the alloy by aging at a temperature of about 800 to 1100 F.

References Cited in the file of this patent UNITED STATES PATENTS 2,754,203 Vordahl July 10, 1956 2,754,204 Jailee et al. July 10, 1956 2,819,194 Herres et a1. Jan. 7, 1958 2,918,367 Crossley et al Dec. 22, 1959 FOREIGN PATENTS 718,822 Germany Mar. 24, 1942 UNITED STATES PATENT OFFICE CERTIFICATE OF CORRECTION Patent No. 3, 147 115 September 1, 1964 Milton Ba Vordahl It is hereby certified that error appears in the above numbered patent requiring correction and that the said Letters Patent should read as corrected below.

Column 1, line 37, for "0,02%" read 002% line 54 for "treat" read treated column 2, line 19, for "this" read thus column 8, TABLE VII in the sub-heading to the talole for P2-130661" read R2-130661 same column TABUE VIII, second column line 4 thereof for "64 hrs, at 900 F, -AV" read 64 hrse at 900 F. -AC column 11, TABLE XI, second column line 1 thereof, for "175 000" read 176,000

Signed and sealed this 30th day of March 1965,

(SEAL) Attcst:

"' ERNEST w. sw1DER EDWARD J. BRENNER Aitesting Officer Commissioner of Patents 

1. A HEAT TREATED AND AGE HARDENED TITANIUM-BASE ALLOY CONSISTING ESSENTIALLY OF ABOUT: 10 TO 16% VANADIUM, 9 TO 13% CHROMIUM, 2 TO 5% ALUMINUM, UP TO 8% MANGANESE, UP TO 4% EACH OF IRON AND TIN, UP TO 0.2% IN TOTAL AMOUNT OF CARBON, OXYGEN AND NITROGEN, BALANCE SUBSTANTIALLY TITANIUM, CHARACTERIZED BY AN ULTIMATE STRENGTH ON THE ORDER OF 190,000-240,000 P.S.I., AN 0.2% OFFSET YIELD STRENGTH ON THE ORDER OF 170,000-220,000 P.S.I. AND A TENSILE ELONGATION OF AT LEAST 1%. 